A novel multi-scale microstructure to address the strength/ductility trade off in high strength steel for fusion reactors

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The current steels suffer irradiation-induced hardening and embrittlement, such that they are not adequate for planned commercial fusion reactors. Producing high strength, ductility and toughness is difficult, because inhibiting deformation to produce strength also reduces the amount of work hardening available, and thereby ductility. Here we solve this dichotomy to introduce a high strength and high ductility RAFM steel, produced by a novel thermomechanical process route. A unique trimodal multiscale microstructure is developed, comprising nanoscale and microscale ferrite, and tempered martensite with low-angle nanograins. Processing induces a high dislocation density, which leads to an extremely high number of nanoscale precipitates and subgrain walls. High strength is attributed to the refinement of the ferrite grain size and the nanograins in the tempered martensite, while the high ductility results from a high mobile dislocation density in the ferrite, the higher proportion of MX carbides, and the trimodal microstructure, which improves ductility without impairing strength. Physical sciences/Materials science Physical sciences/Materials science/Materials for energy and catalysis Figures Figure 1 Figure 2 Figure 3 Figure 4 Figure 5 Text Nuclear fusion energy has long been regarded by many as a potential potent source of non-intermittent, low carbon electricity ( 1 – 3 ). Fusion is attractive due to the abundance of fuel (hydrogen and its isotopes) ( 4 ) and short lifespan of the radioactive waste products ( 5 ). However, the service conditions in fusion reactors are extreme, with components subjected to irradiation, neutron bombardment, exposure to helium and hydrogen, and very high temperatures ( 6 – 10 ). In particular, within the plasma-facing fusion first wall and breeder blanket a significant effort is required to develop structural materials behind the plasma-facing surfaces that can survive such conditions (> 600°C) for realistic plant lifetimes; at least years for breeder blanket modules ( 1 , 11 – 14 ). It is important that these materials can be manufactured at scale for future demonstration and commercial fusion power plants, such as the European DEMO (EU DEMO) or the UK Spherical Tokamak for Energy Production (STEP) programmes ( 15 – 17 ), and multi-tonne conventional production routes are attractive compared to the need to establish new process routes and supply chains, which require significant investment and time. Currently, some of the most promising materials for the breeder blanket are Reduced Activation Ferritic/Martensitic (RAFM) steels, due to their superior thermal conductivity, relatively low thermal expansion and resistance to radiation-induced swelling and helium embrittlement ( 11 , 18 – 23 ). Despite international efforts to develop RAFM steels since the 1980s, and more recently in China, Russia, India and South Korea, the use of current RAFM steels is limited. There are some important aspects that will restrict the use of current RAFM steels; for example, irradiation induces hardening and embrittlement at lower service temperatures (250–350°C) and loss of creep strength and embrittlement at high operating temperatures (550–650°C) ( 24 – 30 ). To address this, developments seek to either achieve fully martensitic structures to avoid phase boundaries and abnormal growth of ferrite grains ( 31 ), or introduce an extremely high number density of nanoscale precipitates for strengthening at high temperature and to absorb irradiation defects, for example ODS-RAFM steel ( 32 , 33 ). However, fully martensitic structures lead to reduced ductility, and irradiation induced effects limits the application temperature to 450–500°C. It is also important to note the production of ODS steels is limited to small quantities and results in enhancing hardening performance at lower service temperatures. Unlike automotive steels, which are designed to either resist deformation (anti-intrusion) or to deform and absorb large amounts of energy in a crash scenario, RAFM steels are not required, nor expected, to plastically deform in-service. Rather, the focus is to resist (micro-) cracking and damage, with better high temperature creep resistance. The high operating temperatures in the fusion reactor can lead to very large thermal stresses which may result in catastrophic material failure in the presence of stress concentrators, e.g. cracks, voids or other features on the phase boundaries. Therefore, it is expected that by improving the room temperature elongation to failure, it will be possible to extend the high temperature service life of RAFM steels and improve their resistance to irradiation-induced embrittlement. Therefore, excessive strain- or irradiation hardening is undesirable, while at the same time, ductility, toughening and the ability to resist cracking, e.g. at notches, is desired. In a simple single-phase polycrystalline material, the onset of dislocation slip will occur in the grains with the highest Schmid factor, which results in load transfer to the surrounding grains and eventually through the yield transition to the propagation of deformation to every grain in the material. Subsequent work hardening can be relatively limited. This is problematic, as once the work hardening rate drops below the yield stress, the material can neck at any geometric imperfection. Therefore, as the yield stress is raised, tensile ductility and the ability to blunt a crack generally drops, which gives the well-known strength-ductility trade-off. This is the reason why a single process, such as work hardening, is not able to increase strength without a penalty to ductility. Thus, a range of strengthening mechanisms are required within a single material, often at different length scales, which operate harmoniously to simultaneously provide high strengthening and ductility ( 34 – 39 ). Here, we extend this concept of a spectrum of deformation scales to RAFM steels. By designing a novel thermomechanical process route we have been able to produce 3 distinct, heterogeneous ferrite/martensite grain size populations allowing the combination of high strength and ductility. The ferrite phase is usually avoided due to the ease with which it coarsens, however, we show that ferrite with a non-uniform grain size can in fact be used to enhance the damage tolerance of the steel. The novel process route induces an extremely high dislocation density throughout the microstructure. During heat treatment the high dislocation density subsequently induces an extremely high number density of nanoscale precipitates, and importantly replaces a significant fraction of the M 23 C 6 by cubic (Ti,V)C intragranular carbides, giving better high temperature stability. This novel RAFM steel microstructure can extend necking, without over-reliance on strain hardening for improvement of the ductility. Processing A novel thermomechanical manufacturing process was developed, shown schematically in Fig. 1 , to provide a multi-scale ferrite/martensite structure, expressly designed to give improved strength and ductility. An RAFM steel with a nominal composition of Fe-0.11C-9Cr-1.1W-0.2V-0.07Ta-0.4Mn-0.25Si-0.01Ti was used, which is based on the composition for Eurofer97, but with the addition of 0.25Si (wt%). Si improves strength and ductility, accelerates strain induced ferrite formation, and is generally known to retard cementite formation on cooling of austenite in bainitic ferrite ( 34 ). After reheating the slab to the soaking temperature and breakdown rolling in the austenitic temperature regime, rolling was performed in 3 stages. In the austenitic temperature regime during Stage 1 (1150 − 1100°C) the steel is unable to fully recrystallise due to the relatively high alloy content. Instead, partial recrystallisation results in a highly deformed unrecrystallised austenite core (γ 1 ), decorated by fine recrystallised austenite grains (γ 2 ) on the grain boundaries ( 37 , 38 ). The steel was then rolled in Stage 2 at 950 − 900°C, just above the austenite-to-ferrite transformation, in order to bring about the Deformation Induced Ferrite Transformation (DIFT) ( 55 – 57 ). This nucleates nanoscale DIFT ferrite grains (α 1 ) on both the γ 1 and γ 2 grain boundaries; the steel can then be quenched, which results in a bimodal microstructure with nanoscale ferrite grains (α 1 ) and martensite (α’ 1 ), and microscale ferrite grains (α 2 ). In Stage 3 the steel was rapidly cooled by spray quenching to a warm intercritical (α + γ) temperature, 850 − 800°C, and then immediately rolled. All the large precursor γ 1 grains, small necklace γ 2 grains, and the DIFT ferrite grains (α 1 ) are co-deformed. At this stage, solute partitioning (particularly Cr) from the larger γ 1 grains to the fine necklace austenite (γ 2 ) and the deformation-induced ferrite (α 1 ), is enhanced through dislocation-facilitated pipe diffusion. When the steel is then quenched after warm rolling, this difference in Cr content and grain size results in transformation of the large unrecrystallised γ 1 grains to martensite α’ 1 , while the smaller necklace γ 2 grains transform to ferrite (α 2 ) (both a higher Cr content and smaller grain size suppress Ms). Both Stage 2 and Stage 3 variants were then normalised at 980°C for 1h followed by quenching to room temperature. The α 1 , α 2 and α´ 1 transforms to austenite; the Stage 2 steel with two grain size modalities, while the Stage 3 has three grain size modalities, namely γ 3 , γ 4 and γ 5 in increasing size. Grain growth is retarded by the precipitation of M 23 C 6 (M = Cr,Fe) carbides on the grain boundaries of γ 3 , γ 4 and γ 5 . Each austenite grain also inherits its previous composition as the normalisation time of 1h is insufficient for diffusion of larger substitutional species such as Cr and Mn. On quenching to room temperature, the larger γ 5 grains transform to martensite (α´ 2 ) and the smaller γ 3 and γ 4 grains transform to ferrite (α 3 and α 4 respectively). The martensite transformation of γ 5 → α´ 2 injected a large density of mobile dislocations into the surrounding α 4 grains, similar to that observed in DP steels ( 58 ). Finally, both steels were aged at 750°C for 1.5h, below the A1 temperature where the austenite-to-ferrite transformation is complete. This ageing heat treatment has the effect on the microstructures of (i) tempering the martensite (α T ´) (i.e. C diffusion but not the movement of substitutional species), (ii) allowing the dislocations within the martensite laths to rearrange themselves into subgrains and (iii) nanoscale precipitate formation. Microstructure Figure 2 shows the microstructures of the steel obtained using both the Stage 2 (Fig. 2 (a-d)) and Stage 3 (Fig. 2 (e-l)) processing routes. In Fig. 2(a, e), the short chains of α 4 necklace grains are observed using Electron Backscattered Diffraction (EBSD). These can be distinguished from the tempered martensite grains as the latter have an abundance of Low Angle Grain Boundaries (LAGBs, red lines) within each grain (high angle grain boundaries are in black). The grain size distributions from the EBSD measurements (above 5µm, with a significant fraction below this, but only resolvable in the TEM) are given in Fig.S4. The Stage 3 structure had a higher proportion of fine grains (< 10µm) than the Stage 2. The ferrite fractions were measured as ~ 43% for the Stage 2 and ~ 33% for the Stage 3. As a comparison, the microstructure obtained after Stage 2 processing is shown in Fig. 2 (a-d). Without Stage 3, the necklace α 4 and tempered martensite α T ´ grains are slightly larger than the microstructure after Stage 3 processing (Fig. 2 (e-h)). Figure 2 (c, g) shows the substructure of the α T ´martensite laths for both stages, composed of subgrains, which also possess a high residual dislocation density. In Fig. 2 (f), a bowed boundary is observed, where the boundary unpinned itself from carbides, due to larger interparticle spacing, but is still being retarded/pinned by adjacent carbide precipitates, only allowing localised movement of the boundary. Furthermore, Fig. 2 (b, f) reveals a number of (Fe,Cr) 23 C 6 carbides, decorating both the ferrite and tempered martensite grain boundaries in both Stage 2 and Stage 3 microstructures. A bright field STEM micrograph and corresponding STEM-EDS images are shown in Figs. 2 (i-l); (Fe,Cr) 23 C 6 carbides are mostly confined to the tempered martensite. In Fig. 2 (d, h), the dislocation structure in the α 4 grains is significantly different between the Stage 2 and Stage 3 processing routes. In the Stage 3 steel α 4 grains, the high dislocation density present after cooling from normalisation rearrange to form Low Energy Dislocation Structures (LEDS) ( 59 , 60 ). These dislocations also facilitate pipe diffusion, forming high number density of nanoscale MX carbides (~ 15nm) (nominally, (Ti,V)C) on and in the immediate vicinity of these dislocations, effectively pinning the LEDS in place after Stage 3 processing (Fig. 2 (h)). This is compared to Stage 2 rolling where only a random distribution of dislocations pinned by sparse nanoscale carbides in Fig. 2 (d) is observed. During ageing, there is also pressure for the α 4 grains to coarsen. The α 4 grain boundaries are highly decorated with either carbides or significantly finer α 3 grains. These pinning particles effectively prevent the grain growth of α 4 grains, but only begin to migrate slightly, forming curved interfaces (Fig. 2 (f)). Small-Angle Neutron Scattering (SANS) experiments have been undertaken for the measurement of the precipitation density, Fig. 3 . Figure 3 (a) shows one-dimensional plots of nuclear scattering intensity versus scattering vector on the trimodal RAFM steel developed here, Eurofer97 and pure iron, respectively. Taking the ratio of magnetic to nuclear scattering (R(q)), we can determine that the trimodal RAFM steel has much lower R(q) (a value of ~ 1) than the baseline Eurofer97 (a value of ~ 2), indicating a lower fraction of Cr 23 C 6 type precipitates larger than 150 nm (calculations in Supplementary Table 2). This is significant, as these grain boundary carbides are held to be deleterious to creep performance ( 61 ). The population of ~ 15 nm diameter cubic (Ti,V)C intragranular carbides was found by SANS to make up 0.16% volume fraction in the Stage 3 steel, shown in Fig. 3 and Fig. S2 and S3, while the fraction was nearly zero in Eurofer 97 steel. These nanoscale carbides would be expected to improve strength, without inhibiting ductility, while potentially providing higher tolerance to neutron damage. Tensile properties The tensile properties of both microstructures are shown in Fig. 1 (b), obtained using full-size ASTM E8 sheet samples ( 62 ). Both stage 2 and stage 3 samples had a higher yield strength than the baseline Eurofer97 RAFM steel. Interestingly, the addition of the stage 3 warm rolling gave higher yield strength, most likely owing to improved precipitation strengthening rather than the final dislocation density per se , as evidenced by the SANS bulk measurement results (Fig. 3 ). Moreover, what was striking was the significantly improved total and post-uniform elongation, which lies well outside the normal “banana” relationship for current RAFM and Dual Phase steels, Fig. 1 (c). In order to investigate the origin of the impressive mechanical properties, the accumulation of damage during tensile testing of the Stage 3 samples, a series of interrupted tensile tests were conducted at engineering strains of 9% (true strain: 0.09), 16% (true strain: 0.145), 38% (true strain: 0.32) and at failure (49% (true strain 0.49)), Figs. 4 , S6. These strains were selected on the basis of the work hardening behaviour, Fig.S6. The arrangement of dislocations into cells occurred at a strain of 9%, Fig. 4 (b-d), pinned by (Ti,V)C precipitates at the cell boundaries. By a strain of 16%, the dislocation density in cell walls increased substantially and these became elongated in the applied stress direction with increasing strain until failure (49% strain). The cell interiors have relatively low dislocation density, but in some there are intense slip bands (Fig. S7). Such evidence of planar slip in ferrite has been reported in austenite-ferrite dual phase steels when grain rotation is inhibited ( 68 , 69 ). In this case it is likely that the intense slip bands are due to the very fine ferrite grain size. With continued increase in strain to 38%, two planar slip systems are observed within the network cell structures, (110)[111] and (112)[111], Fig. S7 (d). The lack of forest hardening therefore indicates that there may be strain softening in the ferrite. However, there may nevertheless be considerable post-uniform elongation as the ferrite phase remains soft and ductile up to fracture. Thus, planar slip and elongation of subgrains, leading to softening, combined with the formation of dislocation cell structures, causing strengthening, significantly improves the mechanical properties to achieve extremely high post-uniform elongation deformation. At a strain of 38% a new microstructural feature was observed, hitherto not reported. New fine scale (< 100nm), strain free, ferrite grains appeared, first in the EBSD-Transmission Kikuchi Diffraction (TKD) map (Fig. 4 (k, i)), with the number density increasing at a strain of 49% in BF-STEM as well as the TKD map (Fig. 4 (p, q)). These strain-free grains formed at the ferrite - tempered martensite interfaces at the regions of highest kernel average misorientation, an indicator of GND density. These features have not been seen before and are most likely associated with the intense dislocation activity in the grain structure with the presence of bowed boundaries (Fig. S8), and are likely to be a feature of the high tensile ductility. Figure 5 shows TKD maps from near the fracture surface (necked region). In Fig. 5 (a-c) voids can be observed at both the ferrite/ferrite and ferrite/tempered martensite boundaries. Voids on grain boundaries are not conventionally associated with positive effects. In dual phase (DP) steels higher volume fraction of voids on the dual phase boundaries leads to shorter post-necking elongation. This occurs as those voids act as stress concentrations and lead to dual phase boundary failure. However, it is interesting to note that quite a few of the voids remained small (~ 100nm), despite the material experiencing post-instability deformation. These small voids along with the new small grains of a similar size were found during the tensile test, Fig. 5 . It is interesting that such voids appeared stable, unlike those observed in DP steels, where voiding at the martensite/ferrite interface leads to failure, which is again another feature of the high tensile ductility observed for the Stage 3 processed material. Discussion The effect of the microstructural features developed by this process on both strength and ductility is significant. An increase in yield strength is classically obtained through a reduction in grain size (and with a reduction in grain size distribution ( 63 )). However, the decreased ductility associated with fine grain sizes is a major limitation. For example, ODS steels have high strength, but poor impact toughness, which is a problem for fusion applications, particularly with irradiation-induced hardening and embrittlement problems. Recent work has shown in several metal systems that a bimodal grain size can improve ductility without significantly impairing strength. The finer grains in the structure impart high strength while the larger grains exhibit high work hardening ability, leading to higher ductility ( 35 , 54 , 64 – 67 ); such behaviour is also verified in our bimodal RAFM steel. In this study, the trimodal structure was developed to further improve both strength and ductility for RAFM steel, and this will be a more suitable microstructure for fusion reactor use than bimodal structures due to its greater hardening resistance and tolerance of defects, delaying embrittlement. For example, the trimodal microstructure RAFM steel achieved a remarkable elongation of 49%, which far exceeds that found in current RAFM steels. In conventional DP steels, voids form at stress concentrations at the martensite/ferrite interface, which ultimately limit ductility. In contrast in the present case, fine voids were observed at these interfaces, but these appeared stable and were not ductility limiting. Moreover, very fine new strain free grains were formed during tensile deformation, which, while not fully explained, are clearly a reflection of stable deformation to high strain. The key features in the microstructure that contributed to such high ductility include the high mobile dislocation density within the ferrite, which, with the absence of forest hardening, indicates that there may be strain softening in the ferrite. This allows considerable post-uniform elongation as the ferrite remains soft and ductile up to fracture. In addition, the higher proportion of fine MC carbides improve strength without greatly inhibiting ductility, and the nanoscale subgrains in martensite and nanoscale ferrite grains also both benefit the strength. Conclusion In summary, a new thermomechanical rolling process has been applied to an RAFM steel composition which produces a completely different microstructure to that conventionally seen in RAFM steel after heat treatment. This microstructure comprises micron-sized ferrite, ferrite with a size in the nanoscale range, and tempered martensite with subgrain structures, combining fine (Ti, V)C precipitates formed on the high density of dislocations, further pinning the structure and adding to the strength, named a trimodal microstructure. The newly designed RAFM steels achieved similar ductility to highly ductile interstitial free (IF) steels, but with substantially higher strength, giving the best combination of strength and ductility. This is a significant breakthrough in the design of RAFM steels, never accomplished before and will provide an RAFM steel with both the desired high temperature strength, sufficiently low impact transition temperature and potentially high tolerance radiation damage. Declarations Acknowledgements This work is supported by the EPSRC grant EP/X030652/1, Royal Society Grant RG\R2\232517 and SUSTAIN Research Hub Early Career Research (application number ECRC1 014 Gong). The authors wish to acknowledge the Engineering and Physical Sciences Research Council (grant number EP/S018107/1) as a part of ‘SUSTAIN Manufacturing Hub’, and Henry Royce Institute for Advanced Materials, funded through EPSRC grants EP/R00661X/1 and EP/P02470X/1 for access to the JEOL JEM-F200 and JEOL JEM 7900F. The authors also wish to acknowledge the DARE project (grant number EP/L025213/1) and Professor Cameron Pleydell-Pearce from Swansea University for their support in hot rolling. The authors also acknowledges the Science and Technology Facilities Council (STFC) for granting access to neutron beamtime at ISIS, ZOOM and Sans2d facilities. The authors also wish to thank the Karlsruhe Institute of Technology for supply of the Eurofer97 for the SANS measurement. This work has been partially funded by the EPSRC Energy Programme (grant number EP/W006839/1), supporting the time of Yiqiang Wang and Huw Dawson. Competing Interests The authors declare that they have no competing interests. Contributions Peng Gong designed the study, performed most of the experimental work, and wrote the manuscript. T.W.J. Kwok undertook melting and rolling of steels, figure construction and discussed the data. R. Goodall discussed the data and contributed to writing the manuscript. Yiqiang Wang undertook SANS. Huw Dawson undertook heat treatment. D. Dye and W. Mark Rainforth contributed to writing the manuscript. Data Availability statement All data needed to evaluate the conclusions in the paper are present in the paper and/or the Supplementary Materials. References A. Morán, R. Coto, J. Belzunce, J. M. Artímez, Experimental development at a pilot plant scale of a Reduced Activation Ferritic/Martensitic RAFM steel. Adv Sci Tech. 73, 36–40 (2010). S. J. Zinkle, L. L. Snead, Designing Radiation Resistance in Materials for Fusion Energy. Nuclear Fusion 57, 92005 (2017). M. Calcagnotto, Y. Adachi, D. Ponge, D. 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Sietsma, Effect of W and N on mechanical properties of reduced activation ferritic/martensitic EUROFER-based steel grades. J. Nucl. Mater. 502, 282–288 (2018). H. He, S. Huang, H. Wang, X. Huang, Isothermal holding processes of a reduced activation ferritic/martensitic steel to form a bainitic/martensitic multiphase microstructure and its mechanical properties. Mater. Sci. Eng. A. 822, 141645 (2021). B. H. Kim, J. Moon, S. D. Kim, J. H. Jang, T. H. Lee, H. U. Hong, H. C. Kim, K. M. Cho, C. H. Lee, Effect of concentrations of Ta and Ti on microstructure and mechanical properties of 9Cr-1W reduced activation ferritic/martensitic steel. Fusion Eng. Des. 151, 111364 (2020). M. Calcagnotto, Y. Adachi, D. Ponge, D. Raabe, Deformation and fracture mechanisms in fine- and ultrafine-grained ferrite/martensite dual-phase steels and the effect of aging. Acta Mater. 59, 658–670 (2011). Z. Liu, D. Li, G. Qiao, Y. Li, Nanometer martensite flakes in high-temperature deformation-induced ferrite grains of a low-carbon steel. ISIJ International 56, 777–780 (2007). Z. Xiong, A. A. Saleh, E. V Pereloma, Strain-induced ferrite formation and its effect on mechanical properties of a dual phase steel produced using laboratory simulated strip casting. J. Alloys Compd. 721, 291-306 (2017) A. Mohamadizadeh, A. Zarei-hanzaki, S. Heshmati-manesh, A. Imandoust, The effect of strain induced ferrite transformation on the microstructural evolutions and mechanical properties of a TRIP-assisted steel. Mater. Sci. Eng. A 607, 621–629 (2014). A. Kundu, D. P. Field, Influence of plastic deformation heterogeneity on development of geometrically necessary dislocation density in dual phase steel. Mater. Sci. Eng. A 667, 435–443 (2016). D. Kuhlmann-Wilsdorf, Theory of Plastic Deformation : -properties of low energy dislocation structures. Mater. Sci. Eng. A 113, 1–41 (1989). C. Laird, P. Charsely, H. Mughrabi, Low Energy Dislocation Structures Produced by Cyclic Deformation. Mater. Sci. Eng. 81, 433–450 (1986) M. Taneike, F. Abe, K. Sawada, Creep-strengthening of steel at high temperatures using nano-sized carbonitride dispersions. Nature . 424, 294–296 (2003). E8/E8M-13a Standard Test Methods for Tension Testing of Metallic Materials, American National Standards Institute. S. Berbenni, V. Favier, M. Berveiller, Impact of the grain size distribution on the yield stress of heterogeneous materials. Int. J. Plast. 23, 114–142 (2007). Y. M. Wang, E. Ma, Three strategies to achieve uniform tensile deformation in a nanostructured metal. Acta Mater. 52, 1699–1709 (2004). Y. Wang, M. Chen, F. Zhou, E. Ma, High tensile ductility in a nanostructured metal. Nature. 419, 912–915 (2002). D. Orlov, J. Zhou, S. Hall, M. Ota-Kawabata, K. Ameyama, Advantages of architectured harmonic structure in structural performance. IOP Conf. Ser. Mater. Sci. Eng. 580, 12019 (2019). Y. Hsieh, L. Zhang, T. Chung, Y. Tsai, J. Yang, In-situ transmission electron microscopy investigation of the deformation behavior of spinodal nanostructured δ-ferrite in a duplex stainless steel. SMM. 125, 44–48 (2016). L. Zhang, R. Song, C. Zhao, F. Yang, Work hardening behavior involving the substructural evolution of an austenite-ferrite Fe-Mn-Al-C steel. Mater. Sci. Eng. A. 640, 225–234 (2015). D. Witkin, Z. Lee, R. Rodriguez, S. Nutt, E. Lavernia, Al-Mg alloy engineered with bimodal grain size for high strength and increased ductility. Scr. Mater. 49, 297–302 (2003). Additional Declarations There is NO Competing Interest. 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Also discoverable on Platform About Our Team In Review Editorial Policies Advisory Board Help Center Resources Author Services Accessibility API Access RSS feed Manage Cookie Preferences © Research Square 2026 | ISSN 2693-5015 (online) Privacy Policy Terms of Service Do Not Sell My Personal Information {"props":{"pageProps":{"initialData":{"identity":"rs-3953989","acceptedTermsAndConditions":true,"allowDirectSubmit":false,"archivedVersions":[],"articleType":"Article","associatedPublications":[],"authors":[{"id":275852169,"identity":"a9453c35-67f2-469d-bab7-6edc15ed1ce2","order_by":0,"name":"William Rainforth","email":"data:image/png;base64,iVBORw0KGgoAAAANSUhEUgAAAZAAAAAyAQMAAABI0h/eAAAABlBMVEX///8AAABVwtN+AAAACXBIWXMAAA7EAAAOxAGVKw4bAAABE0lEQVRIie3RMUsDMRTA8XcEcsu73vpEuX4CIUegOkj9KncUzsWlm1LQAyGT7g4Fv0bHlEC73BeQLr3FqUPdPBA0nrild6tD/kOGhB/JIwA+338tg8SuGGyDEuLwAQKN7T7vItKeIxOWHD0agH4Cv4RTS3QPOS2jel8vRHL28rS6bRYXhCHTegfjIVCRuchIh5LySsiT1aDYRFVByHi2nMMkLanQbsKBcvWZP3McbQJl7i4ZCoPAMqCr8gBhH7kSLZk26sveEu8tue8inP4IREpbgmCJseTAwwzn55ZI4tfyOFKTn1nEci7WqcI39/hrxV4bJRJiVfreqDFhbOrt7mY2jMNCuAgw5y6Ino/0+Xw+X2ffWgVUv/mkVHYAAAAASUVORK5CYII=","orcid":"https://orcid.org/0000-0003-3898-0318","institution":"The University of Sheffield","correspondingAuthor":true,"prefix":"","firstName":"William","middleName":"","lastName":"Rainforth","suffix":""},{"id":275852170,"identity":"77bbc136-28a7-4257-9641-3c39252905d2","order_by":1,"name":"Peng Gong","email":"","orcid":"","institution":"The University of Sheffield","correspondingAuthor":false,"prefix":"","firstName":"Peng","middleName":"","lastName":"Gong","suffix":""},{"id":275852171,"identity":"b8400142-7188-4a9b-b70e-31b0a3eb1525","order_by":2,"name":"Yiqiang Wang","email":"","orcid":"","institution":"United Kingdom Atomic Energy Authority, Culham Science Centre, Abingdon, OX13 3DB","correspondingAuthor":false,"prefix":"","firstName":"Yiqiang","middleName":"","lastName":"Wang","suffix":""},{"id":275852172,"identity":"4c3a9b85-cea3-496b-991d-3d5eb48f31d6","order_by":3,"name":"Thomas Kwok","email":"","orcid":"","institution":"Singapore Institute of Manufacturing Technology","correspondingAuthor":false,"prefix":"","firstName":"Thomas","middleName":"","lastName":"Kwok","suffix":""},{"id":275852173,"identity":"84f3a19a-0ba3-4220-9845-3088e322a145","order_by":4,"name":"Huw Dawson","email":"","orcid":"","institution":"United Kingdom Atomic Energy Authority","correspondingAuthor":false,"prefix":"","firstName":"Huw","middleName":"","lastName":"Dawson","suffix":""},{"id":275852174,"identity":"cb67f1f2-7070-475f-b080-eb9b7fe7cc2a","order_by":5,"name":"Russell Goodall","email":"","orcid":"","institution":"The University of Sheffield","correspondingAuthor":false,"prefix":"","firstName":"Russell","middleName":"","lastName":"Goodall","suffix":""},{"id":275852175,"identity":"f0139a87-b501-4f70-bdb4-8ce4bcc363f0","order_by":6,"name":"David Dye","email":"","orcid":"https://orcid.org/0000-0002-8756-3513","institution":"Imperial College London","correspondingAuthor":false,"prefix":"","firstName":"David","middleName":"","lastName":"Dye","suffix":""}],"badges":[],"createdAt":"2024-02-13 16:51:39","currentVersionCode":1,"declarations":"","doi":"10.21203/rs.3.rs-3953989/v1","doiUrl":"https://doi.org/10.21203/rs.3.rs-3953989/v1","draftVersion":[],"editorialEvents":[{"content":"https://doi.org/10.1038/s41467-025-58042-8","type":"published","date":"2025-03-20T04:00:00+00:00"}],"editorialNote":"","failedWorkflow":false,"files":[{"id":51929331,"identity":"425b33b3-f94f-40ae-a0dc-c9650fe22bb2","added_by":"auto","created_at":"2024-03-04 04:35:31","extension":"png","order_by":1,"title":"Figure 1","display":"","copyAsset":false,"role":"figure","size":910142,"visible":true,"origin":"","legend":"\u003cp\u003e(a) Schematic of the novel thermomechanical manufacturing process and microstructural evolution. Note the bimodal microstructure is obtained without the warm rolling step but also with the same heat treatment. (b) Tensile behaviour of the reference material Eurofer97 and our bimodal- and trimodal- RAFM steels. (c) Comparison of the investigated steel and other literature and commercial Dual Phase (DP) and RAFM steels. The values from (c) for comparison with other high performance materials are reported in literature (\u003cem\u003e40\u003c/em\u003e–\u003cem\u003e54\u003c/em\u003e): Dual Phase steels (40-47); RAFM steels (48-54).\u003c/p\u003e","description":"","filename":"1.png","url":"https://assets-eu.researchsquare.com/files/rs-3953989/v1/fe723c2f089a56bddc1b4754.png"},{"id":51929425,"identity":"3fde07d4-fa9a-4663-bba7-d9fb39134672","added_by":"auto","created_at":"2024-03-04 04:43:31","extension":"png","order_by":2,"title":"Figure 2","display":"","copyAsset":false,"role":"figure","size":1987181,"visible":true,"origin":"","legend":"\u003cp\u003e(a-d) Microstructures of the comparative two-stage bimodal RAFM steel and (e-h) three-stage processing route in the trimodal RAFM steel. (a) EBSD grain boundary map of the bimodal RAFM steel. BF-TEM micrographs of (b) the general microstructure, (c) tempered martensite grains and (d) showing the relatively low dislocation and precipitate density within an α\u003csub\u003e4\u003c/sub\u003e grain. (e) EBSD grain boundary map of the trimodal RAFM steel. Black lines indicate HAGBs (\u0026gt;5˚) and red lines indicate LAGBs (\u0026lt;5˚). BF-TEM microstructures of (f) an α\u003csub\u003e4 \u003c/sub\u003enecklace grain, (g) tempered martensite grains showing a subgrain structure and (h) dislocations pinned by fine precipitates within an α\u003csub\u003e4 \u003c/sub\u003egrain. (i-l) TEM-BF and EDS maps of a ferrite + tempered martensite region in the three-stage processed trimodal RAFM steel, showing the (Fe,Cr)\u003csub\u003e23\u003c/sub\u003eC\u003csub\u003e6 \u003c/sub\u003eand (Ti,V)C carbide locations. In the V map, finer scale (V,Ti)C are also visible (see Fig. S5).\u003c/p\u003e","description":"","filename":"2.png","url":"https://assets-eu.researchsquare.com/files/rs-3953989/v1/8fd94eef6706fbdf3b8218e4.png"},{"id":51929332,"identity":"fb27bd79-47a0-4962-b2cf-bd56bc280d9e","added_by":"auto","created_at":"2024-03-04 04:35:31","extension":"png","order_by":3,"title":"Figure 3","display":"","copyAsset":false,"role":"figure","size":836880,"visible":true,"origin":"","legend":"\u003cp\u003eExperimental configuration used in the current SANS measurements. (a) One-dimensional nuclear scattering intensities versus scattering vector obtained from the Eurofer97, the present trimodal RAFM steel, and pure iron (as a reference). (b) An incident neutron beam transmitted through a bulk specimen containing nano-sized precipitates embedded in a ferritic matrix. (c) The resultant SANS two-dimensional pattern in the presence of a horizontal magnetic field.\u003c/p\u003e","description":"","filename":"3.png","url":"https://assets-eu.researchsquare.com/files/rs-3953989/v1/85e542913797e3d83ab4f846.png"},{"id":51929336,"identity":"bfcd7e2f-da8d-4627-af7a-12814759dbb8","added_by":"auto","created_at":"2024-03-04 04:35:31","extension":"png","order_by":4,"title":"Figure 4","display":"","copyAsset":false,"role":"figure","size":1822055,"visible":true,"origin":"","legend":"\u003cp\u003e(a) Engineering stress-strain curve of the Trimodal RAFM steel. Markers indicate the strain to which interrupted tensile tests were conducted. (b-d) Sample strained to 9%: (b) grain boundary map, red – LAGB (\u0026lt; 5˚), black – HAGB (\u0026gt; 5˚), (c) STEM-BF micrograph of martensite and (d) STEM-BF micrograph of precipitates and dislocations in the ferrite. (e-i) Sample strained to 16%: (e) grain boundary map; (k) Transmission Kikuchi Diffraction (TKD) boundary map and its corresponding KAM in (l); STEM-BF micrographs showing (h) martensite and (i) dislocation cell formation in ferrite. (j-n) Sample strained to 38%: (j) grain boundary map; (k) TKD boundary map with red arrows pointing to the nucleation of low orientation gradient ferrite grains along a ferrite/martensite boundary and corresponding (l) KAM map; STEM-BF micrographs showing (m) martensite and (n) elongated dislocation cell formation in ferrite. (o-s) Sample strained to 49%, \u003cem\u003ei.e. \u003c/em\u003efailure. (o) grain boundary map; (p) TKD boundary map showing periodic formation of strain-free ferrite grains with corresponding (q) KAM map; STEM-BF micrographs of (r) newly formed strain-free grains at a ferrite/martensite boundary and (s) elongated subcell formation in ferrite.\u003c/p\u003e","description":"","filename":"4.png","url":"https://assets-eu.researchsquare.com/files/rs-3953989/v1/3254b8a916022bcd10a4f463.png"},{"id":51929334,"identity":"cdfece2d-9af0-4455-8377-3d0f3d8ee10c","added_by":"auto","created_at":"2024-03-04 04:35:31","extension":"png","order_by":5,"title":"Figure 5","display":"","copyAsset":false,"role":"figure","size":507398,"visible":true,"origin":"","legend":"\u003cp\u003eMicrographs were obtained from a post mortem three-stage trimodal sample just below the fracture surface. (a) TKD image quality map and corresponding (b) phase and boundary map showing several voids along a grain boundary. Black – HAGB, Yellow – LABG, red phase – BCC. (c) TKD image quality map from another region showing similar voiding along grain boundaries.\u003c/p\u003e","description":"","filename":"5.png","url":"https://assets-eu.researchsquare.com/files/rs-3953989/v1/6487746a02c3ec48fd74c2cc.png"},{"id":78942234,"identity":"604be4ca-03fb-416f-8f9b-2846fe1ffb8b","added_by":"auto","created_at":"2025-03-21 07:06:48","extension":"pdf","order_by":0,"title":"","display":"","copyAsset":false,"role":"manuscript-pdf","size":7952197,"visible":true,"origin":"","legend":"","description":"","filename":"manuscript.pdf","url":"https://assets-eu.researchsquare.com/files/rs-3953989/v1/d0730272-6765-46e7-a2a7-60bde08d3f96.pdf"},{"id":51929335,"identity":"04c12185-a1f5-4c72-ae5e-7a0bbeec4223","added_by":"auto","created_at":"2024-03-04 04:35:31","extension":"docx","order_by":1,"title":"","display":"","copyAsset":false,"role":"supplement","size":1502734,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cbr\u003e\u003c/p\u003e","description":"","filename":"Supplementary.docx","url":"https://assets-eu.researchsquare.com/files/rs-3953989/v1/147ab76d5f49abf19a9c8cd2.docx"}],"financialInterests":"There is \u003cb\u003eNO\u003c/b\u003e Competing Interest.","formattedTitle":"A novel multi-scale microstructure to address the strength/ductility trade off in high strength steel for fusion reactors","fulltext":[{"header":"Text","content":"\u003cp\u003eNuclear fusion energy has long been regarded by many as a potential potent source of non-intermittent, low carbon electricity (\u003cspan additionalcitationids=\"CR2\" citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR3\" class=\"CitationRef\"\u003e3\u003c/span\u003e). Fusion is attractive due to the abundance of fuel (hydrogen and its isotopes) (\u003cspan citationid=\"CR4\" class=\"CitationRef\"\u003e4\u003c/span\u003e) and short lifespan of the radioactive waste products (\u003cspan citationid=\"CR5\" class=\"CitationRef\"\u003e5\u003c/span\u003e). However, the service conditions in fusion reactors are extreme, with components subjected to irradiation, neutron bombardment, exposure to helium and hydrogen, and very high temperatures (\u003cspan additionalcitationids=\"CR7 CR8 CR9\" citationid=\"CR6\" class=\"CitationRef\"\u003e6\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR10\" class=\"CitationRef\"\u003e10\u003c/span\u003e). In particular, within the plasma-facing fusion first wall and breeder blanket a significant effort is required to develop structural materials behind the plasma-facing surfaces that can survive such conditions (\u0026gt;\u0026thinsp;600\u0026deg;C) for realistic plant lifetimes; at least years for breeder blanket modules (\u003cspan citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e, \u003cspan additionalcitationids=\"CR12 CR13\" citationid=\"CR11\" class=\"CitationRef\"\u003e11\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR14\" class=\"CitationRef\"\u003e14\u003c/span\u003e). It is important that these materials can be manufactured at scale for future demonstration and commercial fusion power plants, such as the European DEMO (EU DEMO) or the UK Spherical Tokamak for Energy Production (STEP) programmes (\u003cspan additionalcitationids=\"CR16\" citationid=\"CR15\" class=\"CitationRef\"\u003e15\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR17\" class=\"CitationRef\"\u003e17\u003c/span\u003e), and multi-tonne conventional production routes are attractive compared to the need to establish new process routes and supply chains, which require significant investment and time.\u003c/p\u003e \u003cp\u003eCurrently, some of the most promising materials for the breeder blanket are Reduced Activation Ferritic/Martensitic (RAFM) steels, due to their superior thermal conductivity, relatively low thermal expansion and resistance to radiation-induced swelling and helium embrittlement (\u003cspan citationid=\"CR11\" class=\"CitationRef\"\u003e11\u003c/span\u003e, \u003cspan additionalcitationids=\"CR19 CR20 CR21 CR22\" citationid=\"CR18\" class=\"CitationRef\"\u003e18\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR23\" class=\"CitationRef\"\u003e23\u003c/span\u003e). Despite international efforts to develop RAFM steels since the 1980s, and more recently in China, Russia, India and South Korea, the use of current RAFM steels is limited. There are some important aspects that will restrict the use of current RAFM steels; for example, irradiation induces hardening and embrittlement at lower service temperatures (250\u0026ndash;350\u0026deg;C) and loss of creep strength and embrittlement at high operating temperatures (550\u0026ndash;650\u0026deg;C) (\u003cspan additionalcitationids=\"CR25 CR26 CR27 CR28 CR29\" citationid=\"CR24\" class=\"CitationRef\"\u003e24\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR30\" class=\"CitationRef\"\u003e30\u003c/span\u003e). To address this, developments seek to either achieve fully martensitic structures to avoid phase boundaries and abnormal growth of ferrite grains (\u003cspan citationid=\"CR31\" class=\"CitationRef\"\u003e31\u003c/span\u003e), or introduce an extremely high number density of nanoscale precipitates for strengthening at high temperature and to absorb irradiation defects, for example ODS-RAFM steel (\u003cspan citationid=\"CR32\" class=\"CitationRef\"\u003e32\u003c/span\u003e, \u003cspan citationid=\"CR33\" class=\"CitationRef\"\u003e33\u003c/span\u003e). However, fully martensitic structures lead to reduced ductility, and irradiation induced effects limits the application temperature to 450\u0026ndash;500\u0026deg;C. It is also important to note the production of ODS steels is limited to small quantities and results in enhancing hardening performance at lower service temperatures.\u003c/p\u003e \u003cp\u003eUnlike automotive steels, which are designed to either resist deformation (anti-intrusion) or to deform and absorb large amounts of energy in a crash scenario, RAFM steels are not required, nor expected, to plastically deform in-service. Rather, the focus is to resist (micro-) cracking and damage, with better high temperature creep resistance. The high operating temperatures in the fusion reactor can lead to very large thermal stresses which may result in catastrophic material failure in the presence of stress concentrators, \u003cem\u003ee.g.\u003c/em\u003e cracks, voids or other features on the phase boundaries. Therefore, it is expected that by improving the room temperature elongation to failure, it will be possible to extend the high temperature service life of RAFM steels and improve their resistance to irradiation-induced embrittlement. Therefore, excessive strain- or irradiation hardening is undesirable, while at the same time, ductility, toughening and the ability to resist cracking, e.g. at notches, is desired.\u003c/p\u003e \u003cp\u003eIn a simple single-phase polycrystalline material, the onset of dislocation slip will occur in the grains with the highest Schmid factor, which results in load transfer to the surrounding grains and eventually through the yield transition to the propagation of deformation to every grain in the material. Subsequent work hardening can be relatively limited. This is problematic, as once the work hardening rate drops below the yield stress, the material can neck at any geometric imperfection. Therefore, as the yield stress is raised, tensile ductility and the ability to blunt a crack generally drops, which gives the well-known strength-ductility trade-off. This is the reason why a single process, such as work hardening, is not able to increase strength without a penalty to ductility. Thus, a range of strengthening mechanisms are required within a single material, often at different length scales, which operate harmoniously to simultaneously provide high strengthening and ductility (\u003cspan additionalcitationids=\"CR35 CR36 CR37 CR38\" citationid=\"CR34\" class=\"CitationRef\"\u003e34\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR39\" class=\"CitationRef\"\u003e39\u003c/span\u003e).\u003c/p\u003e \u003cp\u003eHere, we extend this concept of a spectrum of deformation scales to RAFM steels. By designing a novel thermomechanical process route we have been able to produce 3 distinct, heterogeneous ferrite/martensite grain size populations allowing the combination of high strength and ductility. The ferrite phase is usually avoided due to the ease with which it coarsens, however, we show that ferrite with a non-uniform grain size can in fact be used to enhance the damage tolerance of the steel. The novel process route induces an extremely high dislocation density throughout the microstructure. During heat treatment the high dislocation density subsequently induces an extremely high number density of nanoscale precipitates, and importantly replaces a significant fraction of the M\u003csub\u003e23\u003c/sub\u003eC\u003csub\u003e6\u003c/sub\u003e by cubic (Ti,V)C intragranular carbides, giving better high temperature stability. This novel RAFM steel microstructure can extend necking, without over-reliance on strain hardening for improvement of the ductility.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e"},{"header":"Processing","content":"\u003cp\u003eA novel thermomechanical manufacturing process was developed, shown schematically in Fig. \u003cspan class=\"InternalRef\"\u003e1\u003c/span\u003e, to provide a multi-scale ferrite/martensite structure, expressly designed to give improved strength and ductility. An RAFM steel with a nominal composition of Fe-0.11C-9Cr-1.1W-0.2V-0.07Ta-0.4Mn-0.25Si-0.01Ti was used, which is based on the composition for Eurofer97, but with the addition of 0.25Si (wt%). Si improves strength and ductility, accelerates strain induced ferrite formation, and is generally known to retard cementite formation on cooling of austenite in bainitic ferrite (\u003cspan class=\"CitationRef\"\u003e34\u003c/span\u003e). After reheating the slab to the soaking temperature and breakdown rolling in the austenitic temperature regime, rolling was performed in 3 stages. In the austenitic temperature regime during Stage 1 (1150\u0026thinsp;\u0026minus;\u0026thinsp;1100\u0026deg;C) the steel is unable to fully recrystallise due to the relatively high alloy content. Instead, partial recrystallisation results in a highly deformed unrecrystallised austenite core (\u0026gamma;\u003csub\u003e1\u003c/sub\u003e), decorated by fine recrystallised austenite grains (\u0026gamma;\u003csub\u003e2\u003c/sub\u003e) on the grain boundaries (\u003cspan class=\"CitationRef\"\u003e37\u003c/span\u003e, \u003cspan class=\"CitationRef\"\u003e38\u003c/span\u003e).\u003c/p\u003e\n\u003cp\u003eThe steel was then rolled in Stage 2 at 950\u0026thinsp;\u0026minus;\u0026thinsp;900\u0026deg;C, just above the austenite-to-ferrite transformation, in order to bring about the Deformation Induced Ferrite Transformation (DIFT) (\u003cspan class=\"CitationRef\"\u003e55\u003c/span\u003e\u0026ndash;\u003cspan class=\"CitationRef\"\u003e57\u003c/span\u003e). This nucleates nanoscale DIFT ferrite grains (\u0026alpha;\u003csub\u003e1\u003c/sub\u003e) on both the \u0026gamma;\u003csub\u003e1\u003c/sub\u003e and \u0026gamma;\u003csub\u003e2\u003c/sub\u003e grain boundaries; the steel can then be quenched, which results in a bimodal microstructure with nanoscale ferrite grains (\u0026alpha;\u003csub\u003e1\u003c/sub\u003e) and martensite (\u0026alpha;\u0026rsquo;\u003csub\u003e1\u003c/sub\u003e), and microscale ferrite grains (\u0026alpha;\u003csub\u003e2\u003c/sub\u003e).\u003c/p\u003e\n\u003cp\u003eIn Stage 3 the steel was rapidly cooled by spray quenching to a warm intercritical (\u0026alpha;\u0026thinsp;+\u0026thinsp;\u0026gamma;) temperature, 850\u0026thinsp;\u0026minus;\u0026thinsp;800\u0026deg;C, and then immediately rolled. All the large precursor \u0026gamma;\u003csub\u003e1\u003c/sub\u003e grains, small necklace \u0026gamma;\u003csub\u003e2\u003c/sub\u003e grains, and the DIFT ferrite grains (\u0026alpha;\u003csub\u003e1\u003c/sub\u003e) are co-deformed. At this stage, solute partitioning (particularly Cr) from the larger \u0026gamma;\u003csub\u003e1\u003c/sub\u003e grains to the fine necklace austenite (\u0026gamma;\u003csub\u003e2\u003c/sub\u003e) and the deformation-induced ferrite (\u0026alpha;\u003csub\u003e1\u003c/sub\u003e), is enhanced through dislocation-facilitated pipe diffusion. When the steel is then quenched after warm rolling, this difference in Cr content and grain size results in transformation of the large unrecrystallised \u0026gamma;\u003csub\u003e1\u003c/sub\u003e grains to martensite \u0026alpha;\u0026rsquo;\u003csub\u003e1\u003c/sub\u003e, while the smaller necklace \u0026gamma;\u003csub\u003e2\u003c/sub\u003e grains transform to ferrite (\u0026alpha;\u003csub\u003e2\u003c/sub\u003e) (both a higher Cr content and smaller grain size suppress Ms).\u003c/p\u003e\n\u003cp\u003eBoth Stage 2 and Stage 3 variants were then normalised at 980\u0026deg;C for 1h followed by quenching to room temperature. The \u0026alpha;\u003csub\u003e1\u003c/sub\u003e, \u0026alpha;\u003csub\u003e2\u003c/sub\u003e and \u0026alpha;\u0026acute;\u003csub\u003e1\u003c/sub\u003e transforms to austenite; the Stage 2 steel with two grain size modalities, while the Stage 3 has three grain size modalities, namely \u0026gamma;\u003csub\u003e3\u003c/sub\u003e, \u0026gamma;\u003csub\u003e4\u003c/sub\u003e and \u0026gamma;\u003csub\u003e5\u003c/sub\u003e in increasing size. Grain growth is retarded by the precipitation of M\u003csub\u003e23\u003c/sub\u003eC\u003csub\u003e6\u003c/sub\u003e (M\u0026thinsp;=\u0026thinsp;Cr,Fe) carbides on the grain boundaries of \u0026gamma;\u003csub\u003e3\u003c/sub\u003e, \u0026gamma;\u003csub\u003e4\u003c/sub\u003e and \u0026gamma;\u003csub\u003e5\u003c/sub\u003e. Each austenite grain also inherits its previous composition as the normalisation time of 1h is insufficient for diffusion of larger substitutional species such as Cr and Mn. On quenching to room temperature, the larger \u0026gamma;\u003csub\u003e5\u003c/sub\u003e grains transform to martensite (\u0026alpha;\u0026acute;\u003csub\u003e2\u003c/sub\u003e) and the smaller \u0026gamma;\u003csub\u003e3\u003c/sub\u003e and \u0026gamma;\u003csub\u003e4\u003c/sub\u003e grains transform to ferrite (\u0026alpha;\u003csub\u003e3\u003c/sub\u003e and \u0026alpha;\u003csub\u003e4\u003c/sub\u003e respectively). The martensite transformation of \u0026gamma;\u003csub\u003e5\u003c/sub\u003e \u0026rarr; \u0026alpha;\u0026acute;\u003csub\u003e2\u003c/sub\u003e injected a large density of mobile dislocations into the surrounding \u0026alpha;\u003csub\u003e4\u003c/sub\u003e grains, similar to that observed in DP steels (\u003cspan class=\"CitationRef\"\u003e58\u003c/span\u003e).\u003c/p\u003e\n\u003cp\u003eFinally, both steels were aged at 750\u0026deg;C for 1.5h, below the A1 temperature where the austenite-to-ferrite transformation is complete. This ageing heat treatment has the effect on the microstructures of (i) tempering the martensite (\u0026alpha;\u003csub\u003eT\u003c/sub\u003e\u0026acute;) (i.e. C diffusion but not the movement of substitutional species), (ii) allowing the dislocations within the martensite laths to rearrange themselves into subgrains and (iii) nanoscale precipitate formation.\u003c/p\u003e"},{"header":"Microstructure","content":"\u003cp\u003eFigure\u0026nbsp;2 shows the microstructures of the steel obtained using both the Stage 2 (Fig.\u0026nbsp;2 (a-d)) and Stage 3 (Fig.\u0026nbsp;2 (e-l)) processing routes. In Fig.\u0026nbsp;2(a, e), the short chains of α\u003csub\u003e4\u003c/sub\u003e necklace grains are observed using Electron Backscattered Diffraction (EBSD). These can be distinguished from the tempered martensite grains as the latter have an abundance of Low Angle Grain Boundaries (LAGBs, red lines) within each grain (high angle grain boundaries are in black). The grain size distributions from the EBSD measurements (above 5\u0026micro;m, with a significant fraction below this, but only resolvable in the TEM) are given in Fig.S4. The Stage 3 structure had a higher proportion of fine grains (\u0026lt;\u0026thinsp;10\u0026micro;m) than the Stage 2. The ferrite fractions were measured as ~\u0026thinsp;43% for the Stage 2 and ~\u0026thinsp;33% for the Stage 3.\u003c/p\u003e \u003cp\u003eAs a comparison, the microstructure obtained after Stage 2 processing is shown in Fig.\u0026nbsp;2 (a-d). Without Stage 3, the necklace α\u003csub\u003e4\u003c/sub\u003e and tempered martensite α\u003csub\u003eT\u003c/sub\u003e\u0026acute; grains are slightly larger than the microstructure after Stage 3 processing (Fig.\u0026nbsp;2 (e-h)). Figure\u0026nbsp;2 (c, g) shows the substructure of the α\u003csub\u003eT\u003c/sub\u003e\u0026acute;martensite laths for both stages, composed of subgrains, which also possess a high residual dislocation density. In Fig.\u0026nbsp;2 (f), a bowed boundary is observed, where the boundary unpinned itself from carbides, due to larger interparticle spacing, but is still being retarded/pinned by adjacent carbide precipitates, only allowing localised movement of the boundary. Furthermore, Fig.\u0026nbsp;2 (b, f) reveals a number of (Fe,Cr)\u003csub\u003e23\u003c/sub\u003eC\u003csub\u003e6\u003c/sub\u003e carbides, decorating both the ferrite and tempered martensite grain boundaries in both Stage 2 and Stage 3 microstructures. A bright field STEM micrograph and corresponding STEM-EDS images are shown in Figs.\u0026nbsp;2 (i-l); (Fe,Cr)\u003csub\u003e23\u003c/sub\u003eC\u003csub\u003e6\u003c/sub\u003e carbides are mostly confined to the tempered martensite.\u003c/p\u003e \u003cp\u003eIn Fig.\u0026nbsp;2 (d, h), the dislocation structure in the α\u003csub\u003e4\u003c/sub\u003e grains is significantly different between the Stage 2 and Stage 3 processing routes. In the Stage 3 steel α\u003csub\u003e4\u003c/sub\u003e grains, the high dislocation density present after cooling from normalisation rearrange to form Low Energy Dislocation Structures (LEDS) (\u003cspan citationid=\"CR59\" class=\"CitationRef\"\u003e59\u003c/span\u003e, \u003cspan citationid=\"CR60\" class=\"CitationRef\"\u003e60\u003c/span\u003e). These dislocations also facilitate pipe diffusion, forming high number density of nanoscale MX carbides (~\u0026thinsp;15nm) (nominally, (Ti,V)C) on and in the immediate vicinity of these dislocations, effectively pinning the LEDS in place after Stage 3 processing (Fig.\u0026nbsp;2 (h)). This is compared to Stage 2 rolling where only a random distribution of dislocations pinned by sparse nanoscale carbides in Fig.\u0026nbsp;2 (d) is observed. During ageing, there is also pressure for the α\u003csub\u003e4\u003c/sub\u003e grains to coarsen. The α\u003csub\u003e4\u003c/sub\u003e grain boundaries are highly decorated with either carbides or significantly finer α\u003csub\u003e3\u003c/sub\u003e grains. These pinning particles effectively prevent the grain growth of α\u003csub\u003e4\u003c/sub\u003e grains, but only begin to migrate slightly, forming curved interfaces (Fig.\u0026nbsp;2 (f)).\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eSmall-Angle Neutron Scattering (SANS) experiments have been undertaken for the measurement of the precipitation density, Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e3\u003c/span\u003e. Figure\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e3\u003c/span\u003e (a) shows one-dimensional plots of nuclear scattering intensity versus scattering vector on the trimodal RAFM steel developed here, Eurofer97 and pure iron, respectively. Taking the ratio of magnetic to nuclear scattering (R(q)), we can determine that the trimodal RAFM steel has much lower R(q) (a value of ~\u0026thinsp;1) than the baseline Eurofer97 (a value of ~\u0026thinsp;2), indicating a lower fraction of Cr\u003csub\u003e23\u003c/sub\u003eC\u003csub\u003e6\u003c/sub\u003e type precipitates larger than 150 nm (calculations in Supplementary Table\u0026nbsp;2). This is significant, as these grain boundary carbides are held to be deleterious to creep performance (\u003cspan citationid=\"CR61\" class=\"CitationRef\"\u003e61\u003c/span\u003e). The population of ~\u0026thinsp;15 nm diameter cubic (Ti,V)C intragranular carbides was found by SANS to make up 0.16% volume fraction in the Stage 3 steel, shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e3\u003c/span\u003e and Fig. S2 and S3, while the fraction was nearly zero in Eurofer 97 steel. These nanoscale carbides would be expected to improve strength, without inhibiting ductility, while potentially providing higher tolerance to neutron damage.\u003c/p\u003e"},{"header":"Tensile properties","content":"\u003cp\u003eThe tensile properties of both microstructures are shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003e(b), obtained using full-size ASTM E8 sheet samples (\u003cspan citationid=\"CR62\" class=\"CitationRef\"\u003e62\u003c/span\u003e). Both stage 2 and stage 3 samples had a higher yield strength than the baseline Eurofer97 RAFM steel. Interestingly, the addition of the stage 3 warm rolling gave higher yield strength, most likely owing to improved precipitation strengthening rather than the final dislocation density \u003cem\u003eper se\u003c/em\u003e, as evidenced by the SANS bulk measurement results (Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e3\u003c/span\u003e). Moreover, what was striking was the significantly improved total and post-uniform elongation, which lies well outside the normal \u0026ldquo;banana\u0026rdquo; relationship for current RAFM and Dual Phase steels, Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003e(c).\u003c/p\u003e \u003cp\u003eIn order to investigate the origin of the impressive mechanical properties, the accumulation of damage during tensile testing of the Stage 3 samples, a series of interrupted tensile tests were conducted at engineering strains of 9% (true strain: 0.09), 16% (true strain: 0.145), 38% (true strain: 0.32) and at failure (49% (true strain 0.49)), Figs.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e4\u003c/span\u003e, S6. These strains were selected on the basis of the work hardening behaviour, Fig.S6. The arrangement of dislocations into cells occurred at a strain of 9%, Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e4\u003c/span\u003e (b-d), pinned by (Ti,V)C precipitates at the cell boundaries. By a strain of 16%, the dislocation density in cell walls increased substantially and these became elongated in the applied stress direction with increasing strain until failure (49% strain). The cell interiors have relatively low dislocation density, but in some there are intense slip bands (Fig. S7). Such evidence of planar slip in ferrite has been reported in austenite-ferrite dual phase steels when grain rotation is inhibited (\u003cspan citationid=\"CR68\" class=\"CitationRef\"\u003e68\u003c/span\u003e, \u003cspan citationid=\"CR69\" class=\"CitationRef\"\u003e69\u003c/span\u003e). In this case it is likely that the intense slip bands are due to the very fine ferrite grain size. With continued increase in strain to 38%, two planar slip systems are observed within the network cell structures, (110)[111] and (112)[111], Fig. S7 (d). The lack of forest hardening therefore indicates that there may be strain softening in the ferrite. However, there may nevertheless be considerable post-uniform elongation as the ferrite phase remains soft and ductile up to fracture. Thus, planar slip and elongation of subgrains, leading to softening, combined with the formation of dislocation cell structures, causing strengthening, significantly improves the mechanical properties to achieve extremely high post-uniform elongation deformation.\u003c/p\u003e \u003cp\u003eAt a strain of 38% a new microstructural feature was observed, hitherto not reported. New fine scale (\u0026lt;\u0026thinsp;100nm), strain free, ferrite grains appeared, first in the EBSD-Transmission Kikuchi Diffraction (TKD) map (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e4\u003c/span\u003e (k, i)), with the number density increasing at a strain of 49% in BF-STEM as well as the TKD map (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e4\u003c/span\u003e (p, q)). These strain-free grains formed at the ferrite - tempered martensite interfaces at the regions of highest kernel average misorientation, an indicator of GND density. These features have not been seen before and are most likely associated with the intense dislocation activity in the grain structure with the presence of bowed boundaries (Fig. S8), and are likely to be a feature of the high tensile ductility.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eFigure\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e5\u003c/span\u003e shows TKD maps from near the fracture surface (necked region). In Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e5\u003c/span\u003e (a-c) voids can be observed at both the ferrite/ferrite and ferrite/tempered martensite boundaries. Voids on grain boundaries are not conventionally associated with positive effects. In dual phase (DP) steels higher volume fraction of voids on the dual phase boundaries leads to shorter post-necking elongation. This occurs as those voids act as stress concentrations and lead to dual phase boundary failure. However, it is interesting to note that quite a few of the voids remained small (~\u0026thinsp;100nm), despite the material experiencing post-instability deformation. These small voids along with the new small grains of a similar size were found during the tensile test, Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e5\u003c/span\u003e. It is interesting that such voids appeared stable, unlike those observed in DP steels, where voiding at the martensite/ferrite interface leads to failure, which is again another feature of the high tensile ductility observed for the Stage 3 processed material.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e"},{"header":"Discussion","content":"\u003cp\u003eThe effect of the microstructural features developed by this process on both strength and ductility is significant. An increase in yield strength is classically obtained through a reduction in grain size (and with a reduction in grain size distribution (\u003cspan citationid=\"CR63\" class=\"CitationRef\"\u003e63\u003c/span\u003e)). However, the decreased ductility associated with fine grain sizes is a major limitation. For example, ODS steels have high strength, but poor impact toughness, which is a problem for fusion applications, particularly with irradiation-induced hardening and embrittlement problems. Recent work has shown in several metal systems that a bimodal grain size can improve ductility without significantly impairing strength. The finer grains in the structure impart high strength while the larger grains exhibit high work hardening ability, leading to higher ductility (\u003cspan citationid=\"CR35\" class=\"CitationRef\"\u003e35\u003c/span\u003e, \u003cspan citationid=\"CR54\" class=\"CitationRef\"\u003e54\u003c/span\u003e, \u003cspan additionalcitationids=\"CR65 CR66\" citationid=\"CR64\" class=\"CitationRef\"\u003e64\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR67\" class=\"CitationRef\"\u003e67\u003c/span\u003e); such behaviour is also verified in our bimodal RAFM steel. In this study, the trimodal structure was developed to further improve both strength and ductility for RAFM steel, and this will be a more suitable microstructure for fusion reactor use than bimodal structures due to its greater hardening resistance and tolerance of defects, delaying embrittlement.\u003c/p\u003e \u003cp\u003eFor example, the trimodal microstructure RAFM steel achieved a remarkable elongation of 49%, which far exceeds that found in current RAFM steels. In conventional DP steels, voids form at stress concentrations at the martensite/ferrite interface, which ultimately limit ductility. In contrast in the present case, fine voids were observed at these interfaces, but these appeared stable and were not ductility limiting. Moreover, very fine new strain free grains were formed during tensile deformation, which, while not fully explained, are clearly a reflection of stable deformation to high strain. The key features in the microstructure that contributed to such high ductility include the high mobile dislocation density within the ferrite, which, with the absence of forest hardening, indicates that there may be strain softening in the ferrite. This allows considerable post-uniform elongation as the ferrite remains soft and ductile up to fracture. In addition, the higher proportion of fine MC carbides improve strength without greatly inhibiting ductility, and the nanoscale subgrains in martensite and nanoscale ferrite grains also both benefit the strength.\u003c/p\u003e"},{"header":"Conclusion","content":"\u003cp\u003eIn summary, a new thermomechanical rolling process has been applied to an RAFM steel composition which produces a completely different microstructure to that conventionally seen in RAFM steel after heat treatment. This microstructure comprises micron-sized ferrite, ferrite with a size in the nanoscale range, and tempered martensite with subgrain structures, combining fine (Ti, V)C precipitates formed on the high density of dislocations, further pinning the structure and adding to the strength, named a trimodal microstructure. The newly designed RAFM steels achieved similar ductility to highly ductile interstitial free (IF) steels, but with substantially higher strength, giving the best combination of strength and ductility. This is a significant breakthrough in the design of RAFM steels, never accomplished before and will provide an RAFM steel with both the desired high temperature strength, sufficiently low impact transition temperature and potentially high tolerance radiation damage.\u003c/p\u003e"},{"header":"Declarations","content":"\u003cp\u003e\u003cstrong\u003eAcknowledgements\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eThis work is supported by the EPSRC grant EP/X030652/1, Royal Society Grant RG\\R2\\232517 and SUSTAIN Research Hub Early Career Research (application number ECRC1 014 Gong). The authors wish to acknowledge the Engineering and Physical Sciences Research Council (grant number EP/S018107/1) as a part of \u0026lsquo;SUSTAIN Manufacturing Hub\u0026rsquo;, and Henry Royce Institute for Advanced Materials, funded through EPSRC grants EP/R00661X/1 and EP/P02470X/1 for access to the JEOL JEM-F200 and JEOL JEM 7900F. The authors also wish to acknowledge the DARE project (grant number EP/L025213/1) and Professor Cameron Pleydell-Pearce from Swansea University for their support in hot rolling. The authors also acknowledges the Science and Technology Facilities Council (STFC) for granting access to neutron beamtime at ISIS, ZOOM and Sans2d facilities. The authors also wish to thank the Karlsruhe Institute of Technology for supply of the Eurofer97 for the SANS measurement. This work has been partially funded by the EPSRC Energy Programme (grant number EP/W006839/1), supporting the time of Yiqiang Wang and Huw Dawson.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eCompeting Interests\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eThe authors declare that they have no competing interests.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eContributions\u0026nbsp;\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003ePeng Gong designed the study, performed most of the experimental work, and wrote the manuscript.\u0026nbsp;T.W.J. Kwok undertook melting and rolling of steels, figure construction and discussed the data.\u0026nbsp;R. Goodall discussed the data and contributed to writing the manuscript. Yiqiang Wang undertook SANS. Huw Dawson undertook heat treatment. D. Dye and W. Mark Rainforth\u0026nbsp;contributed to writing the manuscript.\u0026nbsp;\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eData Availability statement\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eAll data needed to evaluate the conclusions in the paper are present in the paper and/or the Supplementary Materials.\u003c/p\u003e"},{"header":"References","content":"\u003col\u003e\n \u003cli\u003eA. Mor\u0026aacute;n, R. Coto, J. Belzunce, J. M. Art\u0026iacute;mez, Experimental development at a pilot plant scale of a Reduced Activation Ferritic/Martensitic RAFM steel. Adv Sci Tech. 73, 36\u0026ndash;40 (2010).\u003c/li\u003e\n \u003cli\u003eS. J. Zinkle, L. L. Snead, Designing Radiation Resistance in Materials for Fusion Energy. Nuclear Fusion 57, 92005 (2017).\u003c/li\u003e\n \u003cli\u003eM. Calcagnotto, Y. Adachi, D. Ponge, D. Raabe, Deformation and fracture mechanisms in fine- and ultrafine-grained ferrite/martensite dual-phase steels and the effect of aging. 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Mater. 49, 297\u0026ndash;302 (2003).\u003c/li\u003e\n\u003c/ol\u003e"}],"fulltextSource":"","fullText":"","funders":[],"hasAdminPriorityOnWorkflow":false,"hasManuscriptDocX":true,"hasOptedInToPreprint":true,"hasPassedJournalQc":"","hasAnyPriority":true,"hideJournal":false,"highlight":"","institution":"","isAcceptedByJournal":true,"isAuthorSuppliedPdf":false,"isDeskRejected":"","isHiddenFromSearch":false,"isInQc":false,"isInWorkflow":false,"isPdf":false,"isPdfUpToDate":true,"isWithdrawnOrRetracted":false,"journal":{"display":true,"email":"[email protected]","identity":"nature-portfolio","isNatureJournal":true,"hasQc":false,"allowDirectSubmit":false,"externalIdentity":"","sideBox":"","snPcode":"","submissionUrl":"","title":"Nature Portfolio","twitterHandle":"","acdcEnabled":false,"dfaEnabled":false,"editorialSystem":"ejp","reportingPortfolio":"","inReviewEnabled":true,"inReviewRevisionsEnabled":false},"keywords":"","lastPublishedDoi":"10.21203/rs.3.rs-3953989/v1","lastPublishedDoiUrl":"https://doi.org/10.21203/rs.3.rs-3953989/v1","license":{"name":"CC BY 4.0","url":"https://creativecommons.org/licenses/by/4.0/"},"manuscriptAbstract":"\u003cp\u003eAs well as having suitable mechanical performance, fusion reactor materials for the first wall and blanket must be both radiation tolerant and low activation, which has resulted in the development of reduced activation ferritic/martensitic (RAFM) steels. The current steels suffer irradiation-induced hardening and embrittlement, such that they are not adequate for planned commercial fusion reactors. Producing high strength, ductility and toughness\u003cstrong\u003e \u003c/strong\u003eis difficult, because inhibiting deformation to produce strength also reduces the amount of work hardening available, and thereby ductility. Here we solve this dichotomy to introduce a high strength and high ductility RAFM steel, produced by a novel thermomechanical process route. A unique trimodal multiscale microstructure is developed, comprising nanoscale and microscale ferrite, and tempered martensite with low-angle nanograins. Processing induces a high dislocation density, which leads to an extremely high number of nanoscale precipitates and subgrain walls. High strength is attributed to the refinement of the ferrite grain size and the nanograins in the tempered martensite, while the high ductility results from a high mobile dislocation density in the ferrite, the higher proportion of MX carbides, and the trimodal microstructure, which improves ductility without impairing strength.\u003c/p\u003e","manuscriptTitle":"A novel multi-scale microstructure to address the strength/ductility trade off in high strength steel for fusion reactors","msid":"","msnumber":"","nonDraftVersions":[{"code":1,"date":"2024-03-04 04:35:26","doi":"10.21203/rs.3.rs-3953989/v1","editorialEvents":[],"status":"published","journal":{"display":true,"email":"[email protected]","identity":"nature-communications","isNatureJournal":true,"hasQc":false,"allowDirectSubmit":false,"externalIdentity":"NCOMMS","sideBox":"Learn more about [Nature Communications](http://www.nature.com/ncomms/)","snPcode":"","submissionUrl":"https://mts-ncomms.nature.com/","title":"Nature Communications","twitterHandle":"","acdcEnabled":true,"dfaEnabled":true,"editorialSystem":"ejp","reportingPortfolio":"Nature Communications","inReviewEnabled":true,"inReviewRevisionsEnabled":false}}],"origin":"","ownerIdentity":"d046c2e7-db4b-42a9-8a31-5f84d611d8f5","owner":[],"postedDate":"March 4th, 2024","published":true,"recentEditorialEvents":[],"rejectedJournal":[],"revision":"","amendment":"","status":"published-in-journal","subjectAreas":[{"id":29085306,"name":"Physical sciences/Materials science"},{"id":29085307,"name":"Physical sciences/Materials science/Materials for energy and catalysis"}],"tags":[],"updatedAt":"2025-03-21T07:06:40+00:00","versionOfRecord":{"articleIdentity":"rs-3953989","link":"https://doi.org/10.1038/s41467-025-58042-8","journal":{"identity":"nature-communications","isVorOnly":false,"title":"Nature Communications"},"publishedOn":"2025-03-20 04:00:00","publishedOnDateReadable":"March 20th, 2025"},"versionCreatedAt":"2024-03-04 04:35:26","video":"","vorDoi":"10.1038/s41467-025-58042-8","vorDoiUrl":"https://doi.org/10.1038/s41467-025-58042-8","workflowStages":[]},"version":"v1","identity":"rs-3953989","journalConfig":"researchsquare"},"__N_SSP":true},"page":"/article/[identity]/[[...version]]","query":{"redirect":"/article/rs-3953989","identity":"rs-3953989","version":["v1"]},"buildId":"qtupq5eGEP_6zYnWcrvyt","isFallback":false,"isExperimentalCompile":false,"dynamicIds":[84888],"gssp":true,"scriptLoader":[]}